Materials Science & Engineering A Microstructural characterization and mechanical properties of nano-scale/sub-micron TiB-reinforced titanium matrix composites fabricated by laser powder bed fusion --Manuscript Draft-Manuscript Number: MSEA-D-21-02529 Article Type: Research Paper Keywords: laser powder bed fusion; Titanium matrix composite; Ti-6Al-4V; TiB; Strength; Wear resistance Corresponding Author: Eskandar Fereiduni Hamilton, ON CANADA First Author: Eskandar Fereiduni Order of Authors: Eskandar Fereiduni Ali Ghasemi Mohamed Elbestawi Abstract: A B 4 C/Ti-6Al-4V(Ti64) composite powder containing a minor B 4 C content (0.2wt.%) was developed by a novel technique and was subjected to the laser powder bed fusion (L-PBF) process within a wide range of laser powers and scanning speeds to fabricate titanium matrix composite (TMC) parts. The relative density measurement results revealed that almost fully dense TMC parts could be achieved by optimizing the process parameters . Compared to the Ti64 case, slightly higher values were required in the TMC system to achieve the highest relative density. Microstructural characterisations of the TMC parts revealed the formation of large columnar prior β grains containing in-situ formed nano-scale/sub-micron TiB needles homogenously dispersed in a martensitic matrix. While having almost the same ductility, the fabricated TMC parts showed 25, and 8% higher nanohardness, and compressive yield strength, respectively, and 12% lower wear rate than the Ti64 sample. The improved mechanical properties of the TMC part were due to the contribution of several factors including the incorporation of nano-scale/sub-micron TiB reinforcement, the refinement of the martensite αΥ laths, and the solid solution strengthening effects of carbon atoms. The contribution of TiB presence and solid solution strengthening was found to be ~70%, and ~30% in the overall yield strength enhancement of the TMC parts, respectively. Suggested Reviewers: Essa Khamis University of Birmingham School of Mechanical Engineering k.e.a.essa@bham.ac.uk Jill Urbanic University of Windsor jurbanic@uwindsor.ca Tushkar Borkar Cleveland State University T.Borkar@csuohio.edu Powered by Editorial Manager® and ProduXion Manager® from Aries Systems Corporation Cover Letter Dear Professor Lavernia, Please find enclosed our manuscript entitled "Microstructural characterization and mechanical properties of nano-scale/sub-micron TiB-reinforced titanium matrix composites fabricated by laser powder bed fusion", which we would like to submit for publication in your respected journal, Materials Science and Engineering: A. Metal matrix composites (MMCs) are outstanding engineering materials with tailorable properties, having a huge potential to be used in automotive, aerospace, biomedical, and defense industries. Nevertheless, their application has been overshadowed by alloys and superalloys over the past few decades due to the difficulties associated with the processing of MMCs. Laser powder bed fusion (LPBF) process is one of the most promising additive manufacturing (AM) techniques in terms of fabricating MMCs with fine features and high dimensional accuracy, which has attracted a great deal of attention in recent years. Due to the absence of different nozzles for feeding composite constituents, fabrication of MMCs by the LPBF process requires a composite powder feedstock as the starting material. Since no commercial powder is currently available for fabricating MMC parts, different methods have been utilized in recent years to develop composite powder feedstocks for LPBF applications, among which the ball milling is the most frequently used process. Although the ball milling provides the attachment of the reinforcing powder constituent to the metallic powder, it causes the deformation of the metallic powder particles from spherical to non-spherical, and adversely affects the flowability (as a major characteristic of the powder for LPBF process) of the obtained composite powder. In this study, a novel technique is proposed to produce 0.2wt.%B4C/Ti64 composite powders for the LPBF process, which benefits from the advantages of the ball milling and regular mixing methods. The obtained composite powder had a flowability and an apparent packing density close to those of the monolithic Ti64 powder. A wide range of process parameters were utilized to fabricate parts out of the composite and monolithic Ti64 powders. The process parameters leading to the highest densification levels were employed to fabricate compression, and wear test specimens. Results revealed that almost fully dense TMC parts could be achieved by optimizing the process parameters. Microstructural characterisations of the TMC parts revealed the formation of large columnar prior β grains containing in-situ formed nano-scale/sub-micron TiB needles homogenously dispersed in a martensitic matrix. While having almost the same ductility, the fabricated TMC parts showed 25, and 8% higher nanohardness, and compressive yield strength, and 12% lower wear rate than the Ti64 counterpart, respectively. The improved mechanical properties of the TMC part were due to the contribution of several factors including the incorporation of nano-scale/sub-micron TiB reinforcement, the refinement of the martensite αΥ laths, and the solid solution strengthening effects of carbon atoms. While ~70% of this improvement was due to the direct and indirect strengthening effects caused by the TiB phase, the remaining ~30% was originating from the solid solution strengthening mechanism. We confirm that this manuscript has not been published elsewhere and is not under consideration for publication by another journal. I would highly appreciate it if you could consider this manuscript for publication in the Materials Science and Engineering: A journal. Kind regards, Eskandar Fereiduni Department of Mechanical Engineering, McMaster University Hamilton, Ontario, Canada Manuscript Click here to view linked References Microstructural characterization and mechanical properties of nano-scale/submicron TiB-reinforced titanium matrix composites fabricated by laser powder bed fusion Eskandar Fereiduni *, Ali Ghasemi *, Mohamed Elbestawi Department of Mechanical Engineering, McMaster University, Hamilton, ON L8S 4L7, Canada * Corresponding authors: Eskandar Fereiduni (fereidue@mcmaster.ca), and Ali Ghasemi (ghasemia@mcmaster.ca) Abstract A B4C/Ti-6Al-4V(Ti64) composite powder containing a minor B4C content (0.2wt.%) was developed by a novel technique and was subjected to the laser powder bed fusion (L-PBF) process within a wide range of laser powers and scanning speeds to fabricate titanium matrix composite (TMC) parts. The relative density measurement results revealed that almost fully dense TMC parts could be achieved by optimizing the process parameters. Compared to the Ti64 case, slightly higher πΈπ£ values were required in the TMC system to achieve the highest relative density. Microstructural characterisations of the TMC parts revealed the formation of large columnar prior β grains containing in-situ formed nano-scale/sub-micron TiB needles homogenously dispersed in a martensitic matrix. While having almost the same ductility, the fabricated TMC parts showed 25, and 8% higher nanohardness, and compressive yield strength, respectively, and 12% lower wear rate than the Ti64 sample. The improved mechanical properties of the TMC part were due to the contribution of several factors including the incorporation of nano-scale/sub-micron TiB reinforcement, the refinement of the martensite αΥ laths, and the solid solution strengthening effects of carbon atoms. The contribution of TiB presence and solid solution strengthening was found to be ~70%, and ~30% in the overall yield strength enhancement of the TMC parts, respectively. Keywords: Laser powder bed fusion; Titanium matrix composite; Ti-6Al-4V; TiB; Strength; Wear resistance 1. Introduction Ti-6Al-4V (referred to as Ti64 hereafter) is the most widely used titanium alloy with an (α+β) two phase microstructure, which makes of more than half the usage of titanium alloys. This alloy offers a good balance of properties such as high specific strength, adequate stiffness, appropriate high-temperature stability and resistance, good fatigue behavior as well as outstanding corrosion resistance, making it applicable in a wide range of industries such as aerospace, petrochemical, and biomedical [1, 2]. Among these industries, the aerospace sector accounts for >70% of the whole Ti consumption worldwide [3]. Other than these industries, Ti alloys are rarely used in other sectors due to their relatively high cost [4]. Although having numerous benefits, the limited wear resistance of Ti alloys is a major concern when high abrasion and erosion resistance is required [2]. Accordingly, a great deal of attempt has been made in recent decades toward adding alloying elements or incorporating reinforcing particles (ex-situ and in-situ) to/into Ti alloys to improve not only their hardness and strength, but also their wear resistance [5]. In the as-cast condition, the microstructure of Ti64 alloy consists of very coarse prior β grains with sizes in the order of few mm, making it necessary to employ several thermo-mechanical processing steps with the purpose of breaking these coarse grains down to sub-mm or micron length scales [6]. The melting, thermomechanical processing and final machining stages employed to produce finished components of Ti alloys have been reported to be very expensive, accounting for approximately 60% of the total cost [7]. A vast majority of research has been carried out in recent decades to reduce the cost of finished Ti products and expand their usage to other industrial sectors. For this purpose, it has been tried out to decrease the number of processing steps of Ti alloys via adding alloying elements which are capable of refining the grain size of the as-cast microstructure. The refined microstructures will no longer require a number of high-temperature processing steps, leading to a significant drop in the price of manufactured components. Among a wide variety of elements acting as grain refiner in Ti alloys, boron (B) has an incredible grain refining effect and can reduce the as-cast grain size of Ti alloys by an order of magnitude even when existing as minor amounts (e.g., 0.1wt.%) [8]. Although the reduced cost of Ti alloys obtained by the alloy design eliminates the high temperature processing steps, Ti alloy components obtained by conventional manufacturing processes are still expensive. Due to its numerous advantages over traditional manufacturing, additive manufacturing (AM) has become a direct manufacturing technology with applications across a variety of industries [9, 10]. This unprecedented technology provides fabrication of customized and near-net-shape parts with complex shapes, fine features, and unique geometries in one shot, making it economically attractive and of a high interest especially in applications demanding low volume production of relatively expensive materials such as Ti alloys [11-14]. Accordingly, there is a growing global interest in implementing AM technologies providing improved material efficiency and lower cost. Various AM processes have been emerged in recent decades in which the common theme is the consolidation of the feedstock material in a layer-by-layer manner through the localized melting and subsequent solidification, sharing similarities to the cast microstructures. Direct energy deposition (DED) and powder bed fusion (PBF) are the AM categories most frequently used to fabricate metallic parts. The microstructures of as-deposited Ti64 alloy parts obtained by the DED processes are featured by large columnar prior β grains extended over multiple layers, surrounded by grain boundary α extended along the prior β grain boundaries [15, 16]. The grain boundary α is known as the major impediment to the ductility by providing a continuous pathway for the crack propagation [17, 18]. Recently, there has been a great interest toward engineering the microstructure via controlling the process parameters during the AM process. Although this strategy has been found to be successful in tailoring the microstructure within a single component in some alloys (e.g. Inconel 718) [19], the low thermal gradients needed for prior β grains of Ti64 alloy to form in an equiaxed morphology is not feasible in the processing space of the DED and PBF techniques [20]. Also, the relatively narrow range of optimum processing window poses a great challenge to the microstructure tailoring through the engineering of process variables. Therefore, researchers have been seeking alternative ways of grain refinement in the AM-fabricated Ti64 alloy. Successful attempts have been made lately to refine the size of prior β grains in AM-fabricated Ti alloys via trace addition of B [16, 20], LaB6 [21], Si [22], and Be [23]. Due to its significantly high growth restriction factor [24, 25], B has attracted a considerable attention for this purpose. In the DED-fabricated B-modified Ti64 alloys, B has been shown to not only refine both the prior β and α grain size, but also eliminate the grain boundary α phase [16]. However, despite these favorable features, the presence of large TiB needles (> 50 µm) textured along the prior β grain boundaries in the B-modified Ti64 alloy subjected to the wire-arc AM process (with relatively lower cooling rate compared to other AM processes) has been shown to increase the anisotropy in the microstructure and mechanical properties, due to the large TiB needles being highly susceptible to cracking under tensile loading [21]. Since the size of TiB needles is dependent on the cooling rate, AM processes with higher cooling rates (e.g., laser powder bed fusion (LPBF)) can lead to the formation of larger numbers of finer TiB precipitates with more homogeneous distribution, and consequently reduced anisotropy in the microstructure and mechanical properties. B and C elements are known as the most effective grain refiners in Ti alloys. In addition, their presence in Ti alloys can lead to the formation of TiB and TiC phases, which can play a promising role in improving the mechanical properties of Ti alloys [5, 26]. Therefore, minor B4C (as the source of B and C elements) amount of 0.2wt.% was added to the Ti64 powder in this study to produce the composite powder feedstock. The composite powder was produced by a novel mechanical mixing strategy, causing the flowability and apparent packing to be close to those of the monolithic Ti64 powder. A wide range of process parameters were utilized to fabricate parts out of the composite and monolithic Ti64 powders. The process parameters leading to the highest densification levels were employed to fabricate compression, and wear test specimens. Addition of minor B4C to the Ti64 was found to improve the hardness, compressive strengths, and wear resistance. 2. Materials and Experimental Procedure 2.1. Starting materials and preparation of the composite powder feedstock The starting powders used in this research were gas atomized Ti64 alloy and B4C with the nominal chemical compositions reported in [27]. A composite powder feedstock containing 0.2 wt.%B4C (the rest is Ti64) as the starting reinforcing agent was developed using a novel approach benefitting from the advantages of both regular mixing and ball milling processes. Production of each 300 g of composite powder by this technique involved adding 20 g of 2 h-ball milled 3wt.%B4C/Ti64 composite powder to 280 g of monolithic Ti64 powder, followed by regular mixing for 2 h. The mixing of powders was performed using a highperformance planetary Pulverisette 6 machine operating at a fixed rotational speed of 200 rpm. The ballto-powder weight ratio in the ball milling process was set to be 5:1, while the regular mixing was free from balls. The metallic balls added to the system in the ball milling process were made of hardened stainless steel and had a diameter of 10 mm. 2.2. Characterization of the powders The morphology of the starting powders and the produced composite powders were observed using a Vega Tescan scanning electron microscopy (SEM) operating at an accelerating voltage of 20 kV. The optical absorption of Ti64 and B4C powders as well as the developed composite powder was measured using diffuse reflectance spectroscopy (DRS) technique equipped with an UV-Visible-NIR LAMBDA 950 Perkin Elmer spectrophotometer. The integrating sphere had a diameter of 150 mm and was coated with Spectralon with a spectral resolution of 1 nm. A 100% reflectance standard was used as reference to remove the noise. In order to perform the test, each powder sample was placed in a quartz cuvette and sealed prior to mounting on a Teflon sample holder. The light sources were Deuterium (D2) and Tungsten with the wavelength ranges of 200-320 and 320-2500 nm, respectively. Photon Counting photomultiplier tubes (PMT) and Lead Sulfide (PbS), applicable in the wavelength ranges of 200-860.8 and 861-2500 nm, respectively were used as detectors. The flowability and packing density of the Ti64 and composite powders were evaluated by the FT4 Freeman powder rheometer with the procedure thoroughly explained in [27, 28]. 2.3. L-PBF processing An EOS M280 machine (EOS, Krailling, Germany) equipped with a Yb-fiber laser system delivering power levels of up to 400 W was used in this study. An atmosphere of high purity Ar gas was applied to reduce the oxygen content in the build chamber and accordingly minimize the oxidation chance. Cubic parts with the dimensions of 10×10×10 mm3 were printed on a 200 °C-preheated Ti64 build plate (Figure 1(a)). Using a fixed layer thickness (π‘) of 30 µm, and a hatch spacing (β) of 100 µm, varying laser powers (π), and scanning speeds (π) were employed to study the effect of process variables on the quality of the L-PBF fabricated Ti64 and TMC parts (Table 1). By considering the process parameters, the volumetric energy density (πΈπ£ ) is defined as follows [29]: π J πΈπ£ = πβπ‘ [ππ3 ] Eq. 1 Scanning of layers was conducted using a zigzag scanning strategy, alternating 90° between two successive layers. Table 1. The laser powders and scanning speeds employed to fabricate cubic parts. Laser Power, π· [W] Scanning Speed, π [mm/s] π¬π [J/mm3] 100, 150, 200, 250 400, 600, 800, 1000 33-208 Figure 1. Schematic view of the fabricated: (a) cubic samples, (b) wear test specimen, and (c) compression test specimen. As shown in (c), the printed cuboid specimens were subjected to wire electric discharge machining (EDM) to extract cylindrical compression test samples, as per ASTM-E9-09 standard. (Note: The provided dimensions are not proportional to the actual size). The process parameters leading to the highest relative densities in each case (Ti64 and TMC) were employed to print parts for wear and compression tests (Figure 1(b), and (c)). The discs built for the wear test were cut off the plate via wire EDM, ground and then polished according to the standard metallography procedure before testing. 2.4. Microstructural observations Prior to the microstructural characterization, microhardness measurement and nanoindentation test, the cubic parts sectioned through the front plane (Figure 1(a)) were ground and polished according to the standard metallography procedure. The final stage of sample polishing was performed using colloidal silica with an average particle diameter of 0.06 µm. The non-etched sections were observed using a Keyence (Osaka, Japan) VHX digital microscope to compare the defects of the parts qualitatively and quantitatively. The relative density of parts was measured using ImageJ software. The selection of image analysis technique (instead of Archimedes method) for evaluating the density was due to the fact that the bulk density needed for relative density measurements of the TMC samples was unknown. The reported relative density for each sample represents the average of at least 12 measurements. For microstructural studies, the polished sections were chemically etched using Kroll’s Reagent and were observed using a Nikon optical microscopy (OM) as well as a Vega Tescan SEM operating at an accelerated voltage of 10 kV. Electron back-scattered diffraction (EBSD) studies were also carried out to probe the effect of B 4C addition on the grain size, and texture using hardware and software manufactured by FEI. Spatially resolved EBSD maps were acquired at the voltage of 20 keV using a step size of 0.3 µm. 2.5.Mechanical testing The mechanical properties of fabricated parts were evaluated using microhardness and nanohardness measurements, as well as wear, and room temperature compression tests. Microhardness measurements were conducted using a Matsuzawa microhardness testing machine with a load of 500 g and a dwell time of 10 s. The reported microhardness value for each specimen represents the average of at least 5 distinct measurements. Nanohardness measurements were made using Oliver-Pharr method by utilizing an Anton Paar NHT3 nano-indentation tester (Anton Paar, Graz, Austria) equipped with a Berkovich pyramidalshaped indenter tip. Before conducting nanoindentation tests, calibration was performed by using a Fused Silica reference sample. Maximum load of 20 mN, and loading/unloading rate of 40 mN/min were applied on the polished surface of Ti64 and TMC parts in their front view plane. The pause at the maximum load was 10 s for all cases. Compression test was conducted at room temperature using an Instron tensile testing machine as per ASTM E9-19 standard. The strain rate was set to be 0.005 mm/mm/min. The wear performance of samples was evaluated using an Anton Paar standard TRB3 tribometer in accordance with the ASTM G99-95a standard [30] at room temperature. The employed ball was made of alumina, and had a diameter of 6 mm. The applied normal load was 20 N. The linear speed and acquisition rate were set to be 20 cm/s and 100 Hz, respectively. The tests were performed on three different radii of 4, 6, and 8 mm on the top surface of each sample. Three distinct spots of the worn track at each radius of each sample were analyzed with an Alicona microscope to quantify their depth and width and accordingly calculate the wear rate (π) as follows [31]: π π ×π π=πΉ Eq. 2 where π is the wear volume, πΉπ refers to the applied normal force, and π represents the length. 3. Results 3.1. Composite powder feedstock Figure 2(a) shows the micrograph of the 3wt.%B4C/Ti64 composite powder prepared by the ball milling of powder constituents for 2 h. As it is evident, Ti64 powder particles in the produced composite powder were either deformed or almost spherical. The morphological change (spherical to quasi-spherical/flattened) of the deformed Ti64 particles occurred during the ball milling process, in which the ductile Ti64 powder particles experienced severe plastic deformation and micro-forging induced by balls. The ball milling process also provided the attachment of irregular-shape B4C particles to the surface of ductile Ti64 particles, leading to B4C-decorated Ti64 powder particles. The hard B4C particles hitting the surface of Ti64 ones during the ball milling stage also caused the roughening of their surface. The 3wt.%B4C/Ti64 composite powder presented in Figure 2(a) was regularly mixed (balls were absent) with the monolithic Ti64 powder (Figure 2(b)) to develop the desired composite powder feedstock containing 0.2wt.%B4C. Figure 2(c) presents the micrograph of this composite powder, in which most of the Ti64 powder particles were fully spherical and free from B4C particles adhered to their surface since: (i) B4C particles were well bonded to the surface of Ti64 particles of the ball-milled 3wt.%B4C/Ti64 composite powder. Accordingly, they did not find the chance to attach to the Ti64 powder particles added to the system in the regular mixing stage (second stage). (ii) The attachment of B4C particles to the Ti64 powder particles requires an external load which is the impact energy provided by the balls. Due to the absence of balls and consequently the ball-induced impacts in the regular mixing process, this adherence was not facilitated. The variation in the optical absorption of the Ti64 and B4C powders as well as the developed 0.2wt.%B4C/Ti64 composite powder is given in Figure 2(d) as a function of the wavelength. At 1070 nm, which is the wavelength of the L-PBF machine used in this work, Ti64 and B4C powders had optical absorptions of 66 and 83%, respectively. The composite powder showed a laser absorption of 68%, which was slighly higher than that of the Ti64 powder. This can be ascribed to the contribution of two factors: (i) Addition of only a low content (0.2wt.%) of a highly absorptive powder (B4C) to the Ti64 powder to produce the composite powder. Based on the mixture rule principle, this addition enhances the absorption of the composite powder since a fraction of the Ti64 powder constituent is replaced by a higher absorptive material [10, 32]. (ii) The modest change in the morphology (shape and surface roughness) of Ti64 powder particles in the portion of composite powder subjected to the ball milling process. This slight deviation in the morphology acted to increase the frequency of interactions and consequently the optical absorption [33]. Figure 2. SEM micrographs of: (a) 3wt.%B4C/Ti64, (b) monolithic Ti64, and (c) 0.2wt.%B4C/Ti64 powders. (d) Optical absorption versus the wavelength for the Ti64 and B4C powders as well as the developed 0.2wt.%B4C/Ti64 composite powder. The wavelength of 1070 nm represents the laser wavelength of the L-PBF machine used in this study to fabricate parts. The results of the FT4 Freeman powder rheometer test for Ti64 and 0.2wt.%B4C/Ti64 powders are listed in Table 2. As can be seen, incorporation of 0.2wt.%B4C powder to the Ti64 powder led to an increase in the specific energy and adversely affected the flowability (the higher the specific energy, the lower the flowability). Although B4C particles decorated the surface of only a minor fraction of Ti64 powder particles in the developed composite powder, they reduced the flowability by increasing the inter-particle friction and acting as mechanical inter-locking sites during the flow of powder particles [27, 34]. However, it should be noted that since most of Ti64 particles maintained their desired spherical shape, the produced composite powder is believed to possess a superior flowability compared to the same composite powder that could be obtained by a single stage of ball milling for 2 h. The developed composite powder also showed a lower conditioned bulk density (CBD) than that of the monolithic Ti64 powder. This can be attributed to two factors: (i) addition of a less-dense material (B4C) to the Ti64 powder, and (ii) the increased inter-particle friction caused by the presence of irregular-shape B4C particles [27]. Table 2. Flowability and packing density of powders obtained by the FT4 Freeman powder rheometer technique. Flowability Density Specific Energy, SE [mJ/g] Conditioned Bulk Density, CBD [g/mL] Ti64 2.1 ± 0.09 2.7 ± 0.02 0.2wt.%B4C/Ti64 3.07 ± 0.05 2.38 ± 0.01 Powder 3.2. Densification level and processability Figure 3 shows the variation in the density of the L-PBF fabricated Ti64 and TMC parts as a function of the πΈπ£ . In both cases, the density first increased, and then followed a decreasing trend by increasing the πΈπ£ , leading to the maximum densities being achieved within an optimum range of πΈπ£ . The measurements also revealed that a slightly higher πΈπ£ was required in the TMC case to obtain the maximum density (50 J/mm3 for the TMC vs. 42 J/mm3 for the Ti64 system). To gain a better understanding of the effect of process parameters on the defect formation, non-etched cross-sections of the front view are provided in Figure 4 and Figure 5 for Ti64 and TMC samples, respectively. The porosities observed in samples with relatively low πΈπ£ are characterized by their irregular shape (Figure 4(g) and Figure 5(g)), and are known to be caused by the inter-track and/or inter-layer insufficient overlap. At the laser power of 100 W, the decrease in the scanning speed (increase in πΈπ£ ) led to lower levels of porosity in both Ti64 and TMC parts (Figure 4(a), (d) and Figure 5(a), (d)). However, compared to the Ti64 case, the TMC system required a lower scanning speed (higher πΈπ£ ) to yield an almost defect-free part (600 mm/s vs. 800 mm/s in the case of Ti64). Referring to Figure 3, the onset of decreasing trend in the density started earlier in Ti64 compared to the TMC scenario (83 vs. 104 J/mm3). Moreover, a wider range of πΈπ£ led to parts with densities>99% in the TMC system. As can be descerned in Figure 3, at πΈπ£ levels higher than the optimum range for each case, TMCs possessed relative densities higher than those of Ti64 counterparts at any given πΈπ£ (see Figure 4(f), and Figure 5(f)). This suggests that some of the πΈπ£ values which are higher than the optimum value for Ti64, are still within the optimum processing window for the TMC system. The porosities formed at relatively high πΈπ£ levels featured a spherical/semi-spherical shape (Figure 4(c), and Figure 5(c)), giving evidence of the keyhole mode as the defect formation mechanism under this condition [35, 36]. Figure 3. The variation of relative density versus the volumetric energy density (π¬π ) for Ti64 and TMC parts. The designated regions refer to the π¬π range in which highly dense samples were achieved. Figure 4. Non-etched OM images from the front view cross-sections of Ti64 parts fabricated by different sets of laser powders and scanning speeds. Figure 5. Non-etched OM images from the front view cross-sections of TMC parts fabricated by different sets of laser powders and scanning speeds. 3.3. Microstructure evolution and phase analysis Figure 6(a), and (f) show 3D view microstructures of the L-PBF processed Ti64 and TMC samples subjected to the same πΈπ£ of ~83 J/mm3, respectively. Higher magnification micrographs of the selected regions from the front and top view sections of Ti64 and TMC parts are also presented in Figure 6(b, c), and Figure 6(g, h), respectively. Referring to Figure 6(a), and (b), the front and top sections revealed the formation of long (several millimeter) prior β grains extended over multiple layers along the building direction. The measurements showed two different size scales for the width of these prior β grains. Accordingly, they were labeled as “wide” and “narrow” prior β grains with a mean width of 86±5.9 µm, and 9.4±3.1 µm, respectively. The narrow grains were located between the wide ones. Microstructural observation of the top view showed that the wide prior β grains had a cubic section and were surrounded by the narrow ones. Figure 6(d) shows the EBSD orientation map, while Figure 6(e) presents the band contrast image from the front view section of the Ti64 sample. The microstructure was composed of colonies of αΥ martensite laths with different inclinations mainly at ~±45° relative to the building direction. Each colony represents several αΥ laths having the same variant. Based on the Burgers relationship, twelve different α/αΥ variants can be identified for Ti64, accounting for three α/αΥ variants for each of the four <111>β directions [37, 38]. The XRD analysis results of the Ti64 sample (Figure 7(c)) also featured a fully martensitic microstructure, similar to that of the starting Ti64 powder (Figure 7(b)). The formation of a fully martensitic microstructure in the L-PBF processed Ti64 alloy is due to noticeably high cooling rates of the L-PBF process which well exceeds the critical cooling rate for the martensitic phase transformation of this alloy (410 °C/s) [39, 40]. Referring to Figure 6(d), the wide and narrow prior β grains were distinguishable since αΥ laths transformed from each narrow prior β grain featured a different orientation and color pattern from those of the neighbor wide grains. Referring to front and side views of the TMC sample provided in Figure 6(f), and (g), typical semi-elliptical shape tracks were perceptible due to the employed scanning strategy alternating 90° between subsequent layers. It is worth noting that although the same etchant was used for both Ti64 and TMC samples, the consolidated tracks were not visible in the Ti64 part. Microstructural characterization of the TMC specimen also revealed the formation of elongated alternate wide and narrow prior β grains (Figure 6(f), (g), and (h)). However, compared to the Ti64 sample, narrow prior β grains showed a larger width in the TMC case, and were located along the center of the tracks (Figure 6(g)). The mean width for the wide and narrow prior β grains was measured to be 75.5±6.4 µm, and 19.6±4.1 µm, respectively. The sum of the average width of the wide and narrow prior β grains was almost equal to the employed hatch spacing (100 µm) in both Ti64 and TMC cases. However, the narrow prior β grains surrounding the wide ones were more distinctive and had a greater width in the TMC sample compared to the Ti64 one. This can be ascribed to the slight difference between the heat input of the two samples caused by the difference in the optical absorption and the effective powder layer thickness of the powder systems. Referring to the top views presented in Figure 6, a chessboard pattern of wide square prior β grains surrounded by narrow β grains along their boundaries was obtained. The same chessboard pattern has been previously observed in several research studies dealing with the L-PBF processing of the Ti64 alloy. The formation of this pattern has been attributed to the employed scanning strategy in which the scan directions rotated 90° between subsequent layers [41, 42]. Similar to the Ti64 sample, a fully martensitic matrix was obtained in the L-PBF fabricated TMC system (Figure 6(i), (j)). However, the measurements gave evidence of finer αΥ laths in the TMC part. The mean αΥ lath width was 1.15±0.15, and 0.92±0.12 µm, while the mean length was 15.6±3.1 and 8.4±7.3 µm for the Ti64 and TMC samples, respectively. The size of the α/αΥ laths is an important microstructural feature which governs the mechanical properties of Ti alloys by defining the effective slip length. Smaller α/αΥ colonies generally bring about improved mechanical properties [43, 44]. The XRD phase identification of the TMC part indicated no peaks corresponding to the B4C phase (Figure 7(a), (d)). Given the fact that microstructure of L-PBF fabricated TMC was also free from B4C (Figure 6), it can be concluded that these particles have experienced complete melting/dissolution during the process. The XRD analysis of the TMC part also revealed a fully martensitic microstructure, the same as that of the Ti64 sample (Figure 7(c), (d)). In addition, no peaks corresponding to the TiB and TiC phases were detected. Figure 8 shows SEM micrographs of the TMC part in which fine features are perceptible in the whole microstructure. Due to their needle-like morphology, these precipitates can be characterized as the TiB phase formed through the reaction between B and Ti elements. Higher magnification micrographs of the track interiors (Figure 8(c)) and boundaries (Figure 8(d)) revealed that the fine needle-shape TiB precipitates were slightly coarser and had a larger spacing in the boundaries than interiors. This microstructural difference is believed to be the main reason behind observing the tracks in the TMC part, as opposed to the Ti64 counterpart. Within the track interior, TiB needles were mainly oriented toward the center of the melt pool (Figure 8(c)). This directionality is known to be due the unidirectional heat flux occurring from center toward the boundaries of the melt pool [59, 73, 74], allowing preferential growth of these needle-shape features opposite to the heat flux direction. It is noteworthy that the absence of TiB phase in the XRD pattern of the TMC part (Figure 7) might be due to its noticeably small size. Figure 6. (a) OM 3D view of the microstructure of the L-PBF fabricated Ti64; (b), (c) higher magnification OM images of the selected areas of the front and top views shown in (a), respectively; (d) inverse pole figure (IPF-Z), and (e) band contrast EBSD maps of the front view of the Ti64 sample; (f) OM 3D view of the microstructure of the LPBF fabricated TMC part; (g), (h) higher magnification OM images of the selected areas of the front and top views shown in (f), respectively; (i) inverse pole figure (IPF-Z), and (j) band contrast EBSD maps of the front view of the TMC sample. The employed laser powder, and scanning speed for both Ti64 and TMC samples were 250 W, and 1000 mm/s, respectively. Figure 7. XRD patterns of the starting: (a) B4C, and (b) Ti64 powders, as well as the L-PBF fabricated: (c) Ti64, and (d) TMC parts. Both Ti64 and TMC samples were manufactured using a laser power of 250 W and a scanning speed of 1000 mm/s. Figure 8. Secondary electron (SE)-SEM images from the front view of the L-PBF fabricated TMC part. (a) General view in which the cross-section of a track has been highlighted. Higher magnification micrograph of the selected dashed square in (a) is shown in (b). The square regions marked as “A” and “B” in (b) are shown in a higher magnification in (c) and (d), respectively. 3.4. Mechanical Properties To gain insights into the level of improvement in the mechanical properties caused by the composite fabrication, the hardness, strength, ductility, and wear performance of L-PBF fabricated TMC was evaluated and compared to those of the Ti64 counterpart. As shown in Figure 9, the developed TMC showed 5% higher microhardness than that of the Ti64 part. However, nanoindentation measurements revealed a greater difference between the nanohardness values (25%). It is worth noting that compared to the Ti64 part, a smaller standard deviation was detectable in the nanohardness measurements of the TMC sample. This can be attributed to the enhanced microstructure homogeneity caused by the uniform distribution of fine TiB precipitates in the matrix of the TMC sample. The typical compressive stress-strain curves of the Ti64 and TMC parts are presented in Figure 9(b). Table 3 summarizes the most important properties of these two materials obtained from several tests performed on each material. As it is evident, both samples had almost the same ultimate strength (ππ’ ) and fracture strain (ππ ). However, the TMC sample possessed ~8% higher yield strength (ππ¦ ) than the Ti64 counterpart. Therefore, although the incorporation of nano-scale/sub-micron TiB reinforcement into the Ti64 matrix brought about higher hardness, and strength, it did not have an adverse effect on the ductility. Figure 9. (a) The microhardness and nanohardness values, and (b) typical compressive engineering stress-strain curves of the L-PBF fabricated Ti64 and TMC parts. Table 3. Mechanical properties of the L-PBF fabricated Ti64 and TMC parts obtained by compression testing. Sample ππ¦ ππ’ Young’s Modulus, E [GPa] Fracture Strain, ππ [%] Ti64 TMC 1190 1284 1833 1850 114.4 119.3 21 20.6 Figure 10(a), and (b) shows the variation in the coefficient of friction (COF) for the L-PBF fabricated optimum Ti64 and TMC parts along the sliding distance, respectively. The listed in Table 4, there was no significant difference between the COF for the Ti64 and TMC samples. However, the analysis of the wear tested samples revealed a noticable difference between the depth and width of the worn tracks left on the surface of Ti64 and TMC parts. Figure 10(c), and (d) illustrates representative depth profiles along the width of the worn tracks for Ti64 and TMC samples across the directions shown in Figure 10(d), and (e), respectively. As observed, compared to the Ti64 sample, the TMC featured a shallower depth and a narrower width. The wear rate of samples was calculated by measuring the dimensions of the worn tracks at several spots of each radii and consequently calculating the wear volume. Referring to in Table 4, the TMC sample showed 15% lower wear rate than the Ti64 one. This reveals the remarkable influence of incorporated TiB reinforcements on the improved wear resistance of the L-PBF fabricated TMC. Figure 10. (a), (b) the variation of COF versus sliding distance for Ti64 and TMC samples, respectively. Representative depth-width profile of the worn tracks for (c) Ti64, and (d) TMC samples obtained along the paths depicted in the 3D surface topography images in (e), and (f), respectively. Table 4. The coefficient of friction (COF) and wear rate of L-PBF processed Ti64 and TMC samples. System COF Wear rate [mm3/(N.m)] Ti64 0.4218 1.318×10-3 TMC 0.4314 1.115×10-3 Figure 11 shows the SEM micrographs of the wear tracks for Ti64 and TMC samples. The wear tracks of both samples presented wear scars along the sliding direction which is typical of the wear surfaces. In addition, ploughing grooves were observed in both cases. As can be seen in the higher magnification micrographs in Figure 11(b), and (d), the wear surface contained delamination cracks in the Ti64 and TMC samples. These cracks are generated as a cause of the high strain levels experienced by the surface of the sample when being in contact with the pin during sliding. As indicated in Figure 11(b), and (d), such high strain levels can bring about the delamination of the material during the wear test. A close examination of the wear tracks at several spots of each radii (4, 5, and 8 mm) revealed that compared to the Ti64 sample, the delamination cracks and delamination layers were less intense in the TMC sample. This can be due to the higher hardness and plastic deformation resistivity of the TMC sample caused by the strengthening effects of TiB reinforcements. It has been shown that owing to its high modulus and hardness, incorporation of slight amount of TiB into the Ti alloy matrix can lead to a significant increase in the wear resistance [26]. Figure 11. SEM images of the wear tracks after the sliding wear test for the L-PBF fabricated: (a), (b) Ti64; and (c), (d) TMC samples. Higher magnification micrographs of the selected regions in (a), and (c) and presented in (b), and (d), respectively. 4. Discussions 4.1. Processability As the density measurement results and microstructural observations revealed, addition of a minor amount of B4C (0.2wt.%) to the Ti64 powder changed its L-PBF processability (Figure 3, Figure 4, and Figure 5). This clearly reveals that not only the process parameters but also the powder characteristics and laserpowder interactions need to be re-examined when dealing with a composite powder system containing even minor amounts of a second constituent. As proposed in a very recent study on L-PBF processing of VC/H13 system [45], the volumetric energy density calculation equation most frequently used in L-PBF process needs to be revisited for composite powder systems by taking the laser absorptivity (π΄) and relative powder bed packing density (ππππ ) of the powder also into account as follows: ππ£ = (π΄ × ππππ )πΈπ£ Eq. 3 where ππ£ represents the modified volumetric energy density, and ππππ is defined as πΆπ΅π·/πππ’ππ (πππ’ππ signifies the bulk density of the material). This equation suggests that powders with higher (π΄ × ππππ ) require lower πΈπ£ to yield the optimum ππ£ . Referring to the CBD measurements results provided in Table 2, optical absorption values presented in Figure 2(a), and assuming πππ’ππ to be 4.43 g/cm3 for both Ti64 and TMC bulk samples, the (π΄ × ππππ ) term for the Ti64 and 0.2wt.%B4C/Ti64 systems is 0.4 and 0.36, respectively. This justifies the need for slightly higher πΈπ£ in the composite system (compared to the Ti64) to reach the highest densification level. The observed increasing and then decreasing variation in the density of the Ti64 and TMC samples (Figure 3) is in agreement with the results reported in the literature for the L-PBF processing of metallic materials [46, 47]. The decline in the relative density at relatively high πΈπ£ values exceeding the optimum value has been well documented in the literature and has been shown to be due to the transition of melting mode from conduction to keyhole [29, 35, 48, 49]. During the keyhole mode of melting, the material evaporation caused by application of high πΈπ£ leads to the development of a vapor cavity which further enhances the laser absorption and enables the laser beam to penetrate to a deeper depth. The collapse of the cavity can leave voids in the material which are known as the keyhole porosities and are spherical in shape [48, 50]. As it is perceptible in Figure 3, the decreasing trend in the relative density occurred at a slightly lower πΈπ£ in the Ti64 system compared to the TMC case (83 vs. 104 J/mm3). One of the most effective estimations for the criteria of keyhole threshold has been shown to be based on the normalized enthalpy ( π₯π» ) βπ as follows [48]: π₯π» βπ = π΄π√π· πππ √ππ£π 3 ≥ πππ ππ Eq. 4 in which π₯π» is the specific enthalpy, βπ is the enthalpy at melting defined as πππ , ππ π· is the melting point, ππ is the boiling point, π is the laser beam size, π is the thermal conductivity, and π· is the thermal diffusivity. Eq. 4 takes into account the effect of the applied π, π£, and π on the threshold for the formation of keyhole melting mode. However, it fails to consider the influence of layer thickness as one of the important parameters playing role on the melting mode. Eq. 5 provides the modified threshold for keyhole mode of melting by taking the effective layer thickness also into account [51, 52]. π₯π» βπ 0.865π΄ππ· π √π ππ‘πππ π£ = ππ ≥ πππ ππ Eq. 5 Eq. 5 can be also rewritten as: π π£ ≥ ππππ π‘πππ √π3 Eq. 6 0.865π΄π· Assuming the same π, and ππ for the Ti64 and TMC systems, and considering π‘πππ = π π‘ πππ π π£ 1 ∝ π΄×π Eq. 7 πππ Referring to Eq. 7, the critical [34]: π π£ for the keyhole threshold in the composite system is expected to be 10% higher than that of the Ti64 case. However, based on the experimental results, this difference was ~25%. Although a portion of this discrepancy may have raised because of assuming the same π, and ππ for both systems, there are still other factors related to the composite systems which have not been considered in the currently existing keyhole threshold models. For instance, given the fact that the dissolution of B4C particles during the L-PBF processing of the composite system is an endothermic reaction, a part of the energy absorbed by the powder bed is consumed by their dissolution. 4.2. Microstructure evolution As shown in Figure 6, the microstructure of both Ti64 and TMC parts contained elongated columnar prior β grains which were extended along the building direction. The formation of such columnar prior β grains is a major characteristic of the L-PBF processed Ti64 alloy parts caused by the combination of epitaxial and competitive growth during the solidification. The epitaxial growth deals with the solidification of the melt pool through the growth (rather than nucleation) of β crystals on the β grains of the adjacent track/underlying layer. Those prior β grains having their easy-growth crystallographic direction (<001> direction in cubic crystal structures) perpendicular to the fusion line (along the heat flux direction) experience a preferential growth, meaning that faster growth of these grains crowds out the β grains oriented in less favorable directions [53-55]. As a result, the final solidification structure consists of columnar prior β grains extended along the building direction. As the laser beam interacts with the 0.2wt.B4C/Ti64 composite powder bed, Ti64 powder particles experience melting since their melting temperature is well below the temperature levels induced during the L-PBF process. However, owing to their noticeably high melting point of 3036 K [56] as well as the extremely fast cooling rates associated with the L-PBF process (104-106 K/s), complete melting of B4C particles is rather difficult or even impossible. As the microstructural characterizations (Figure 6 and Figure 8) and XRD analysis results (Figure 7) revealed, the fabricated TMC did not contain B4C constituent in the final microstructure. In fact, B4C particles fully reacted with the surrounding melt. This reaction is believed to occur through the dissolution of B4C particles rather than their melting. The full dissolution of B4C particles in the surrounding molten material is an interesting phenomenon which may not be expected to take place due to the extremely high cooling rate of the L-PBF process and consequently the noticeably short time available for diffusion-based dissolution mechanism. The fast and complete dissolution of highmelting-point B4C particles in this study can be attributed to: (i) their relatively low fraction and small size, and (ii) convective flows (e.g., Marangoni flow, etc.) active within the melt pool during the L-PBF process which consecutively provide fresh molten material (not enriched from B atoms) adjacent to the solid B4C particles [40]. Based on the calculations, the molten material formed by complete dissolution of B4C constituent during the L-PBF processing of the 0.2wt.%B4C/Ti64 composite powder contains 0.156 and 0.039 wt.% B and C, respectively. Figure 12(a) shows the equilibrium phase diagram of the Ti64-B system where the isopleth for the Ti640.156wt.%B system is specified. The schematic illustration of the solidification sequence of this isopleth is also depicted in Figure 12(e)-(h), in which the solidification starts with the formation of β phase from the liquid and their subsequent growth (Figure 12(e), and (f)). Due to the extremely low/negligible solubility of B and C in the β phase, they are rejected from the growing solidification interface upon cooling between the liquidus and eutectic temperatures (Figure 12(a), (f), and (g)), leading to the progressive enrichment of the liquid from B and C. Since the C in the investigated TMC system is noticeably lower than the equilibrium solubility limit of this element in Ti (0.039 vs. 0.125 wt.%), the C content of the enriched liquid is still below this limit, causing C atoms to exist as solid solution in the Ti64 matrix. At the eutectic temperature, the remaining liquid enriched from B transforms to eutectic πππ΅+ π½ (πππ΅πΈ +π½πΈ ) through a pseudo-eutectic reaction (Figure 12 (a), (g), and (h)). Therefore, the equilibrium solidification sequence can be summarized as: πΏ → π½ + πΏ → π½ + (πππ΅πΈ + π½πΈ ) The β phase transforms to (α+β) at the β transus temperature. Therefore, equilibrium microstructure at room temperature consists of lamellar (α+β) islands surrounded by narrow band of (πππ΅πΈ + π½πΈ ). The isopleth of the Ti64 alloy is presented in Figure 12(b), where the solidification starts with the nucleation of β phase from the liquid followed by their growth. Despite the Ti64-0.156wt.%B system with a relatively wide solidification range in the equilibrium state, that of the Ti64 alloy is noticeably narrow. This suggests that no significant partitioning is occurred in the liquid during the solidification of the Ti64 alloy. Figure 12. (a) Equilibrium phase analysis diagram of the Ti64-B system with the B content in the range of 0-0.3 wt.%. The provided isopleth represents the investigated 0.2wt.%B4C/Ti64 system which contains 0.156wt.%B. (b) Equilibrium phase analysis diagram of the Ti-6Al-XV, in which the isopleth of the Ti-6Al-4V (Ti64) alloy has been shown. Scheil-Gulliver solidifications models of the: (c) Ti64-0.2wt.%B4C, and (d) Ti64 systems. The inset in (c) shows the variation of the TiB fraction within the specified temperature range (1530-1420 °C). (e)-(h) Schematic illustration of the solidification sequence of the 0.2wt.%B4C/Ti64 system in the equilibrium condition. The liquid (L) becomes enriched from B and C as the solidification progresses. As the microstructural observations revealed, the fabricated TMC contained fine nano-scale/sub-micron TiB needles homogeneously dispersed in the martensitic matrix. This microstructure differs from the microstructure predicted by the equilibrium phase diagram in Figure 12(a), where the large prior β grains are surrounded by the pseudo-eutectic reaction products (πππ΅πΈ +π½πΈ ) along their boundaries (Figure 12(h)). Although the formation of a martensitic matrix (rather than the (α+β) two-phase microstructure) in the LPBF fabricated TMC can be justified based on the extremely fast cooling rates of the process, the reason behind the deviation of the solidification path from the equilibrium condition needs further analysis. Figure 12(c) illustrates the solidification sequence proposed by the non-equilibrium Scheil-Gulliver model for the TMC system investigated in this study. This model accounts for the infinitely fast diffusion in the liquid and no diffusion in the solid. As can be seen, the Scheil-Gulliver model predicts the same solidification sequence as that of the equilibrium condition, revealing that even the non-equilibrium ScheilGulliver solidification model fails to predict the solidification sequence of the TMC system subjected to the L-PBF process. However, unlike the equilibrium condition in which the TiB phase is formed at a specific temperature (Eutectic temperature of 1440 °C), the Scheil-Gulliver model predicts the formation of TiB to occur within a relatively wide range starting at 1530 °C and finishing at 1440 °C. This can be ascribed to the lack of solute diffusion in the solid based on the Scheil-Gulliver model which provides the liquid with the B content required for the formation of TiB at higher temperatures. Figure 13(a), and (b) shows the microstructure of the last solidified layer of the fabricated TMC part, which is composed of the TiB needles homogeneously dispersed in a martensitic matrix. The microstructure evolution of the last layer is schematically illustrated in Figure 13(c)-(g). Although B has a negligible solubility in the equilibrium condition in the high-temperature β phase (a maximum of ~0.05 wt.% (Figure 13(a))), the extremely fast cooling rate associated with the L-PBF process can significantly diminish the rejection of B atoms into the liquid during the solidification (Figure 13(d), and (e)). For instance, rapid solidification of the Ti-B systems has been shown to cause supersaturation of B in the matrix up to 10 at.% (2.5 wt.%B) [57], which is noticeably higher than the B content of the TMC system in this study (0.156 wt.%). Accordingly, the high-temperature prior β phase is supersaturated from B. Due to the decreased solubility limit during the cooling stage, a portion of the supersaturated B tends to precipitate out as the TiB phase from both the β and the later transformed B-supersaturated martensite (M) phases (Figure 13(f), and (g)). Compared to the bulk of the TMC sample shown in Figure 8, lower volume fraction of TiB phase was perceptible in the last layer. This clearly reveals the significant role of the subsequent thermal cycles (caused by the heat conduction from the melt pool to the previously deposited tracks or layers) on the precipitation of TiB from the B-supersaturated matrix. The same phenomenon has been also observed in the powder metallurgy-processed 1.7wt.%B/Ti64 alloy, where the exposure of the material to elevated temperatures during the thermo-mechanical processing led to the increase in the volume fraction of the TiB phase in the microstructure [58]. Both the equilibrium and Scheil-Gulliver non-equilibrium solidification models are based on the scenario where the temperature gradient is low enough to allow the formation of a constitutionally undercooled zone in front of the solidification interface, favoring the nucleation of new prior β grains. However, when an extremely steep temperature gradient exists (e.g., such as that of the L-PBF process), very limited or even no nucleation events would occur ahead of the solidification interface. This explains the observation that the TMC part also featured columnar prior β grains with almost the same size as those of the Ti64 alloy (Figure 6). It is worth noting that the inability of the minor B amounts in avoiding the formation of columnar prior β grains has been also reported in B-modified Ti64 alloys subjected to the AM processes with thermal gradients noticeably lower than that of the L-PBF process (WAAM) [59]. Figure 13. (a), and (b) SEM micrographs of the last layer of the L-PBF fabricated TMC part. Higher magnification micrograph of the selected region in (a) is shown in (b). (c)-(g) Schematic illustration of the microstructure evolution of the last layer of the L-PBF fabricated TMC part during cooling. “M” in (g) refers to martensite. As observed in Figure 8 and Figure 13, TiB reinforcements had a needle-like shape and were uniformly dispersed in the martensitic matrix. The reason behind such a morphology has been traced back to the strong preferential growth along the [010]27 direction compared to the other directions [60]. The average length and width of TiB needles were 1.6 µm and 260 nm, respectively, being significantly smaller than those formed in Ti-B systems with trace B content subjected to the DED [61], direct laser deposition (DLD) [62], and wire-arc AM (WAAM) [21] processes, which usually have a length in the order of several microns. This corresponds to the noticeably higher cooling rates associated with the L-PBF process. The relatively large TiB needles have been reported to be detrimental to the mechanical properties since they can act as crack nucleation sites during loading [16, 21, 61]. Therefore, the L-PBF fabricated TiB reinforced TMCs are more favorable since they benefit from greater number of finer TiB reinforcements with smaller interreinforcement spacing uniformly dispersed in the matrix (as opposed to fewer numbers of TiB reinforcements with a larger spacing). 4.3. Mechanical Properties Based on the hardness measurements, compression test results, and wear test, the TMC sample benefitted from a higher hardness, yield strength, and wear resistance than the Ti64 counterpart. The improvement in the mechanical properties of the TMC over the Ti64 can be ascribed to the contribution of several strengthening mechanisms, including the grain refinement, second constituent, dislocation, and solid solution. In the following, the level of enhancement in the yield strength caused by each mechanism is discussed. Grain refinement: The increased fraction of boundaries obtained by grain refinement elevates the plastic deformation resistivity and consequently improves the strength of the material [63, 64]. Based on the HallPetch relationship, the increase in the yield strength of the TMC caused by the grain refinement can be expressed as follows [63]: βππ»π = ππ»π ( 1 √π2 − 1 √π1 ) Eq. 8 in which ππ»π is the Hall-Petch coefficient for Ti64; while π2 , and π1 are the mean grain sizes of the TMC and Ti64 samples, respectively. Eq. 8 is typically used for the equiaxed microstructures whereas the Ti64 and TMC samples featured martensitic microstructures characterized by αΥ laths. Therefore, the average width of αΥ laths in each material was taken into account [65]. Second constituent strengthening: The strength improvement caused by the presence of TiB reinforcements in the TMC part is defined as [66]: 1 π βππππ΅ = 2 ππ¦π × ππππ΅ × ππππ΅ π0 Eq. 9 πππ΅ where ππ¦π is the yield strength of the Ti64 matrix, ππππ΅ is the volume fraction of TiB, ( ππππ΅ ) ππππ΅ refers to the aspect ratio of TiB (treated as cylinders with the average length and width of π πππ΅ , and ππππ΅ , respectively), and π0 is the orientation factor which has a value between 0 and 1. Due to the random distribution of TiB whiskers in the matrix of the TMC part, π0 is equal to 0.125 or 0.27 based on the 3D or 2D random array model, respectively [67]. Dislocation strengthening: The increment in the strength originated from the dislocation strengthening can be estimated as follows [68]: 2 βππππ = √(βππππ )2 + (βππ‘βπ )2 + (βππππ ) Eq. 10 in which βππππ is the strength improvement by the Orowan mechanism, βππ‘βπ represents the stress enhancement caused by the thermal expansion mismatch between the TiB and the Ti64 matrix, and βππππ signifies the stress increment due to the geometrically necessary dislocations (GNDs). In the composites containing reinforcements not exceeding 1 µm in size (e.g. TiB needles with nanoscale/submicron width in this study), the reinforcements can act as strong obstacles to the movement of dislocations and impede their motion on the matrix’s slip plane [69]. The strengthening effect caused by this phenomenon is known as the Orowan mechanism, in which the shear component of the applied stress bows out the dislocation between reinforcements. This produces a back stress which in turn prevents the migration of further dislocations and consequently enhances the yield strength [70]. The Orowan stress can be described by the Orowan-Ashby equation as [71]: βππππ = 0.13 πΊπ π΅π π· ln 2ππππ΅ π π Eq. 11 where πΊπ is the shear modulus of the Ti64 matrix, π΅π is the Burgers vector of the Ti64 matrix, π·πππ΅ is the equivalent diameter of the TiB reinforcement, and π is the interparticle spacing. π·πππ΅ and π can be calculated as [68, 72]: 3 π·πππ΅ = √1.5π2πππ΅ π πππ΅ Eq. 12 1 π ≈ π·πππ΅ [(2ππππ΅ )−3 − 1] Eq. 13 Since the thermal expansion coefficients of TiB and Ti64 are very close to each other, βππ‘βπ can be neglected. βππππ deals with the strength improvement caused by the GNDs required to accommodate the plastic deformation mismatch between the TiB reinforcement and the matrix, and is expressed as follows [73]: ππ−π ππ ππππ΅ π·πππ΅ βππππ = 0.4πΊπ √ Eq. 14 in which ππ−π is the compressive true strain of the matrix at the yield point. Solid solution strengthening: As the microstructural observations and phase analysis results revealed, the TiC phase was not detected in the fabricated TMC part. Since the C content in the studied TMC is less than the solid solubility limit of C in Ti, C atoms have been entrapped in the matrix as solid solution. Given the fact that C is an interstitial alloying element, it can generate a significantly strong obstacle (compared to the substitutional elements) for the movement of dislocations in the matrix and consequently lead to higher levels of solid solution strengthening [74]. The solid solution strengthening caused by the interstitial solutes is generally proportional to the square root of the concentration as [75]: βππ π = 0.002πΊπ π1/2 Eq. 15 where π is the carbon concentration in at.%. By taking all the strengthening mechanisms into account, Ramakrishnan’s approach (Eq. 16) can be used to calculate the yield strength of the TMC component (ππ¦−πππΆ ) [76]: ππ¦−πππΆ = ππ¦−π (1 + βππ»π βπ βππππ ) (1 + πππ΅ ) (1 + ) (1 ππ¦−πππΆ ππ¦−πππΆ ππ¦−πππΆ + βππ π ) ππ¦−πππΆ Eq. 16 Eq. 16 estimates the yield strength of the TMC sample to be ~1307 MPa, which is in a good agreement with that obtained by the experiment (1284 MPa). Table 6 summarizes the contribution of each strengthening mechanism in the yield strength increment of the TMC sample over the Ti64 counterpart. While the βππππ΅ is the direct strengthening mechanism originating from the incorporated TiB reinforcement, βππ»π , and βππππ are the indirect strengthening mechanisms activated due to the presence of this phase in the microstructure. Therefore, taking both direct and indirect strengthening mechanisms into consideration, ~70% of the total yield strength enhancement in the TMC is caused by the presence of TiB phase. On the other hand, the remaining ~30% is linked to the solid solution strengthening of the matrix induced by C atoms. The improved hardness and wear resistance of the TMC parts can be also attributed to the same strengthening mechanisms discussed above for the strength. Table 5. The properties and parameters used in obtaining the contribution of different strengthening mechanisms. Properties and Parameters ππ»π π2 π1 ππ πΊπ π0 ππππ΅ π πππ΅ π πππ΅ ππ−π π·πππ΅ Value 300 MPa.µm1/2 0.92 µm 1.15 µm 0.295 nm 45 GPa 0.125 0.05 1590 nm 264 nm 0.0122 0.55 µm Source [77] Measured Measured [78] [79] [67] Measured Measured Measured Measured Measured Table 6. The calculated values of strengthening mechanisms in the TMC part along with their contribution in the enhancement of the yield strength over the Ti64 counterpart. Strengthening Mechanism Value [MPa] Proportion in the total βπ [%] βππ»π 33 29.1 βππππ΅ 22.4 19.8 βππππ 21.3 18.8 βππ π 36.6 32.3 5. Conclusions In this study, the laser powder bed fusion (L-PBF) process was implemented to fabricate high performance titanium matrix composite (TMC) parts using a composite powder feedstock produced by a novel technique. The main conclusions of this research study can be summarized as follows: 1- The proposed composite powder preparation technique benefiting from the advantages of both regular mixing and ball milling was found to yield a 0.2wt.%B4C/Ti64 composite powder having the flowability and apparent packing density close to those of the monolithic Ti64 powder. 2- In both Ti64 and TMC cases, an increasing and then decreasing trend in the relative density of parts was achieved by enhancing the volumetric energy density. Compared to the Ti64 case, slightly higher energy densities were needed in the TMC to achieve almost fully dense samples. Moreover, the decline in the relative density took place at higher energy densities in TMC than Ti64. 3- Microstructural characterizations revealed the formation of large columnar prior β grains in both Ti64 and TMC samples which were elongated along the building direction. The TMC sample featured nano-scale/sub-micron TiB needles homogeneously dispersed in a fully martensitic matrix. 4- The equilibrium and Scheil-Gulliver non-equilibrium solidification models were found not to be able to predict the microstructure evolution experienced by the L-PBF fabricated TMC parts. This was attributed to the extremely high cooling rates of the L-PBF process which led to the supersaturation of B atoms in the β grains during solidification. 5- Although showing almost the same ultimate strength and fracture strain as those of the Ti64 counterpart, the TMC part possessed 5% higher microhardness, 8% higher yield strength, and 12% lower wear rate. 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